High-stability nickel base alloy

ABSTRACT

An improved nickel base alloy is provided long time stability with respect to high rupture and tensile strength and adequate ductility at high temperatures through the avoidance of sigma phase. This is accomplished through the careful balance and coordination of Ti, Al and Cr in a Ni base alloy including, in addition, Co, Mo, Zr, V, B and C.

Ross et al.

[151 3,642,469 [451 Feb. 15, 11972 HIGH-STABILITY NICKEL BASE ALLOY lnventors:

Earl Warren Ross; Howard Thomas Mcllenry, both of Cincinnati, Ohio Assignee:

Filed:

Appl. No.: 859,240

General Electric Company Aug. 29, 1969 Related US. Application Data Continuation-impart of Ser. No. 745,717, July 18, 1968, abandoned, which is a continuation-in-part of Ser. No. 362,945, Apr. 27, 1964, abandoned.

References Cited UNITED STATES PATENTS 3,061,426 10/1962 Bieber ..75/ l 7! Primary Examiner-Richard 0. Dean Attorney-Derek P. Lawrence, Lee H. Sachs, E. S. Lee, Ill, Frank L. Neuhauser, Oscar B. Waddell and Joseph B. Forman [5 7] ABSTRACT An improved nickel base alloy is provided long time stability with respect to high rupture and tensile strength and adequate ductility at high temperatures through the avoidance of sigma phase. This is accomplished through the careful balance and coordination of Ti, Al and Cr in a Ni base alloy including, in

US. Cl ..75/171, 148/325 addition, Co, Mo Zr v, B and Q Int. Cl ..C22e 19/00 Field of Search ..75/l71, 170; 148/32, 32.5 3 Claims, 1 Drawing Figure [Iii/A jiff'fi /fi HIGH-STABILITY NICKEL BASE ALLOY This application is a continuation-in-part of application Serial No. 745,717, filed July 18, 1968, now abandoned which was a continuation-in-part of application Serial No. 362,945, filed Apr. 27, 1964, now abandoned all of common inventors and assignee.

This invention relates to nickel base alloys and, more particularly to castable nickel base alloys of greatly improved stress rupture strength and stability for longtime use.

Some of the most important nickel base alloys in use today are of the type precipitation hardened principally by aluminum and titanium and further strengthened by molybdenum and tungsten. This type of nickel base alloy is particularly useful when made into components for jet engines. One of the most critical applications with regard to temperature and stress in which high strength, ductility and long time stability of the properties during service are required, is as a cast turbine blading material in advanced gas turbine engines. One known alloy, referred to in Bieber US. Pat. No. 3,061,426, has a nominal composition, by weight, of 15% Co, 10% Cr, 5% Ti, 5.5% A1, 3% Mo, 1% V, 0.18% C, 0.015% B, 0.05% Zr with the balance. essentially nickel and incidental impurities. Although this known alloy has good strength and adequate ductility for relatively short times up to about 1,900 F., it has been found that for exposure at 1,400-l,800 F. for times of in excess of 100 hours, and particularly after 300 hours, the alloy structure tends to become unstable and embrittled, hence the ductility and rupture life is reduced substantially. For example, tensile tests at l,400 F resulted in ductility approaching after service exposure of 1,000 hours at 1,600" F. Long time stability tests to 2,500 hours showed that this alloy can become embrittled at l,400 F. after 1,650 F. exposure with or without stress. Thus this alloy, through useful for other purposes, is unsuitable for long time applications as turbine blading material in advanced gas turbine engines.

A principal object of the present invention is to provide an improved nickel base alloy of the above type which has long time stability as evidenced by high rupture and tensile strength and adequate ductility after long exposures under high temperature and stress conditions.

Another object is to provide a cast nickel base alloy, the composition balance of which will suppress the formation of detrimental phases during long time use.

These and other objects and advantages will be more fully understood from the following detailed description and examples which are meant to be exemplary of, rather than limitations on the scope of the present invention.

The drawing is a graphic comparison of the loss of expected life due to excessive sigma phase formation.

Briefly, the present invention involves suppressing particularly in the cast form of nickel base alloys the formation of an undesirable intermetallic phase, with a crystal structure technically referred to as a sigma phase, after long time exposure to elevated temperatures. This is accomplished by lowering or specifically limiting the quantities of elements which tend to bring about formation of the sigma phase. The formation of sigma phase is particularly suppressed through the specific control of the titanium content within the range of 4.0-4.3 weight percent with a careful balance of the other elements which cause phase precipitation. In one more specific form in which many production heats have been made avoiding sigma phase, the alloy of the present invention consists essentially of, by weight, 0.l0.2% C, 8-1 1% Cr, 4.04.3% Ti, -6% Al, 0.03-0.09% Zr, 13-17% Co, 2-4% Mo, 0.7-1.2% V, 0.0l-0.02% B, with the balance nickel and incidental impurities such as Fe, Si, Mn, S, Cu, etc., with the specific condition, however, that the electron vacancy or phase computation number N which will be discussed in detail later, shall not exceed 2.47.

A more complete understanding of the present invention will evolve from a discussion of the detrimental phases which can be present in the above identified known alloy and how such phase is determined, These phases, which are particularly critical in the cast form of the alloy, are greatly suppressed or are entirely eliminated according to the present invention, to result in a different kind of alloy. ln casting the above identified known alloy, a structure called primary gamma prime" solidifies from the liquidus as large spherulitic particles in dendritic areas. This structure is of the Ni (Al, Ti) type which also contains gamma phase and probably fine carbides. These areas containing primary gamma prime have a different chemical composition than other areas in the matrix due to chemical segregation. It has been indicated that they also contain high chromium content which results in reduced available solubility for gamma prime with the probability that more than normal'primary gamma prime for the total Al and Ti content forms at the liquidus in theseareas.

On exposure to elevated temperatures in the known alloy, the gamma and the carbides, which are in the primary gamma prime, agglomerate. For example, these reactions take place at temperatures up to the incipient melting temperature of the primary gamma prime (about 2,250 F At temperatures from l,300-l,800 F sigma plates form in matrix areas surrounding the primary gamma prime. This formation, which is accelerated by stress, appears to be due to the following reac' tions: (1) excessive chromium in the primary gamma prime and surrounding matrix areas first forms grain boundary M C carbides, utilizing carbon from TiC globule carbide breakdown, the primary gamma prime and the matrix, (2) when all of the available carbon is fully tied up as M C it appears that the excessive chromium combines with cobalt, etc., to form a Cr-Co-Mo type sigma.

Long time stability tests show the above identified known alloy can be brittle at 1,400F. and below after 1,400 l ,800" F. long time exposure with or without stress. This loss of ductility has been correlated with the presence of acicular shaped sigma phase in the microstructure as well as the presence of the undesirable M C type carbide in the grain boundary formed after breakdown of MC carbide, predominantly TiC. However, equally important with the effect on tensile ductility, the presence of sigma phase is particularly significant in seriously reducing stress rupture life above 1,400 F.

When the effect of sigma phase is virtually eliminated as a result of the balanced alloy compositions of the present invention, most practically accomplished by reducing the titanium content, a different kind of alloy structure is formed and the reduction in long time rupture properties is eliminated This is accomplished without a sacrifice in other properties.

It has been found specifically that the primary gamma prime content can be considerably reduced with 40-43 weight percent titanium in the presence 'of about 5.5% Al, the balance of the composition being preferably 0.l5-O.20% C, 8 l 1% Cr, 13-17% C0, 0.030.09% Zr, 24% Mo, 0.7-1 .2% V, 0.01-0.02% B with the balance essentially nickel and incidental impurities. Although sigma phase can be eliminated entirely in such alloy range with less than 4.0% Ti, for example at 3.5% Ti, very long time alloy strength properties tend to be reduced.

The formation of sigma phase in cast alloys is dependent upon chemistry, casting segregation, grain size and exposure temperature, time and stress. One feature of the present invention is to provide a chemistry which virtually eliminates the fon'nation of sigma phase by forming a different kind of alloy to greatly improve rupture strength and ductility while maintaining good tensile properties.

Extensive studies have shown that formation of the embrittling sigma phase is accelerated by the presence of stress in the temperature range of l,300-l ,800 F. and particularly at 1,500-l,650 F. The lowest room temperature and 1,400 F. residual tensile ductility were accompanied by the largest quantities of sigma phase. Also, the reduced ductility was due to the normal formation of the grain boundary M C, type carbide where M"is predominantly Cr.

Another significant effect of excessive sigma formation was found to be a loss of rupture life with an accompanying increase in rupture ductility. Long time rupture testing showed that the 1,500 to 1,650 E. rupture lives of sigma prone heats were reduced by as much as 90% from the expected life based on parametric extrapolation of short time rupture properties. It has been found that profuse sigma formation along octahedral planes resulted in intersigmatic rupture fracture from the above identified alloy test bar, with 5.3% Ti, that had failed in 469 hours at l,500 F./55,000 psi. The anticipated life, based on short time properties and without sigma formation during this test would have been 2,000 hours. The rupture results of this sigma prone heat is shown in FIG. 1 compared with the average stress rupture properties of that same alloy on shorter time basis, for example, 100 hours. The loss of rupture life at 1,500" and l,650 F. is easily recognizable. The l,500 F./40,000 p.s.i. test point shown failed in 967 hours versus an expected life of 8,000 hours.

The terminal nature of sigma phase has been discussed by Boesch and Slaney in Metals Progress, July 1964, pages 109 -l l 1. Although sigma phase may be removed by heat treatment, it will recur when the alloy experiences the same time and temperature conditions under which sigma was formed originally. Therefore, a careful composition balance is required to inhibit original sigma phase formation.

In the preliminary evaluation of the alloy of the present invention, comparison was made between the known alloy and certain development heats. The composition of the preliminary heats are shown in the following Table I.

TABLE I Compositions in Weight Base: 0. 1 80. 19 C; 0012-0014 B; 0.07-0.08 Zr;

Bal. Ni and Incidental Impurities In the above Table I, alloy A is the known type alloy, the loss in expected stress rupture life due to excessive sigma formation for which is shown in the drawing. Alloy B was found to be a sigma-free heat 4.0 weight percent titanium. Alloy C was a heat having occasional sigma phase at a 4.5% titanium level.

In accelerated high stress testing for gas turbine airfoils, these initial long time stress rupture properties of the alloys of Examples B and C were considered to be excellent. For example, the l,500 F./50,000 p.s.i. lives have reached over 2,000 hours in six tests compared to the 469 hours for the sigma prone heat of the known alloy of Example A and previously referred to in the drawing. Normal stress levels in such parts are about 15,000 psi.

Although in these accelerated tests the long time rupture lives and residual ductility of the 4.0 and 4.5% titanium level alloys of Examples B and C were considered to be excellent, there was evidence, confirmed by photomicrographic studies of these alloys, of sigma formation in a few of the 4.5% titanium long-time rupture specimens. As was indicated before, the presence of sigma phase, particularly as it affects long time stress rupture properties, is a terminal condition in the type of nickel base alloy to which this invention relates. It cannot be removed permanently by heat treatment.

The short time mechanical properties of these types of alloys, for example at 100 hours, might be of academic interest. However, in their intended application in jet engines, articles made from such alloys must operate reliably for long periods of time without requiring disassembly of complex engines and replacement of component parts, the alloy life of which is depleted. Many thousands of hours, for example 10,000 or more hours of life, are goals of some jet engine components today. Thus when one talks of utility of high-temperature superalloys in such gas turbine art, from a practical viewpoint, such utility must include consideration of long time properties at the temperatures at which those types of alloys are intended to operate. Therefore, the alloy of the present invention recognizes the danger of the presence of sigma phase. In order to eliminate any possibility of sigma formation, this alloy limits the titanium content at a safe 4.3% along with the other ranges specified.

Another means of determining potential sigma formation is through the electron vacancy theory or phase computation identified by the number N in Table I. This type of phase computation, sometimes referred to as PhaComp, is now one of the metallurgical tools used in the metals producing industry to control compositions. One of the early reports on the theory was by Linus Pauling in The Nature of Interatomic Forces in Metals", Physical Review Vol. 54, Dec. 1, 1938. There Pauling related the bonding between atoms of certain elements to the average number of electron vacancy (N,.) in their bonding orbitals. Since that time others have reported various formulas and relationships which produce numbers dependent upon the precipitation assumed to occur in the nickel base alloy. One such method which is sometimes referred to as the Woodyatt-Simms-Beattie method was reported upon in A Contemporary View of Nickel Base Superalloys By C. Sims in Journal of Metals, Oct. I966, pages ll 19 to 1130. In the Tables of this specification, the number resulting from this type ofcalculation is identified by N,.;,.

In the calculation of the N electron vacancy or phase computation number, assumptions are made (1) that all of the boron forms an M B boride; (2) that one-half of the carbon forms MC where M is Ti and the other half of the carbon forms M C where M is predominantly Cr along with Mo and W; and (3) that all of the aluminum, the remaining titanium, 3% of the original chromium and 50% of the vanadium shall combine to form the gamma prime composition Ni (AL, Ti, 0.03 Cr, 0.5 V).

After converting the alloy composition to atomic percent, the residual matrix composition remaining after gamma prime precipitation is calculated. Then the average electron vacancy number N,, is calculated from the following formula:

N, =i m, (N,,), in which:

N is the average electron vacancy number iis each individual element in turn m,is the atomic fraction of each element in the matrix (N,.),-is the electron vacancy number of each respective element Thus this method is one for determining the average electron vacancy concentration per atom remaining in the gamma matrix of a nickel base alloy taking certain reactions into consideration. A low N number heat is less likely to produce embrittling and weakening intermetallic phases such as sigma in the matrix. The particular method (N used in the description of the present invention assumes that the 3% Cr is included with Al and Ti in the gamma prime, and that carbide and boride reactions are accounted for in the calculations.

With reference now to Table I, it can be seen that N of the known alloy A is at 2.93 and that the borderline alloy C which had a slight indication of sigma was at 2.62 compared to 2.47 for the sigma-free alloy of Example B. Therefore, as a further limitation of the alloy of the present invention consistent with evaluation data, the alloy of the present invention is further defined as having a phase computation number N at a maximum of 2.47 coupled with the specified preferred titanium range of 404.3% in one form of the present alloy.

During subsequent evaluation of the alloy of the present invention, a large number of alloy examples were studied, typical examples of which are shown in Table II.

TABLE II Compositions in Weight Base: 0. l30.20% C; 0.0l-0.02% B; 0.03-0.07% Zr Bal. Ni and Incidental Impurities Alloy Cr Co V Mo Al Ti N The compositions of Table II were melted as 30-pound master melts. Cast-to-size bars were obtained by remelting stock from the master melts. All melting was done under vacuum and a fine grain size of 1/6 inch or smaller was specified.

Because the normal critical sigma phase formation range is at about 1,650 F., and the alloy loses rupture strength with increasing sigma phase formation, a stress rupture test in air was developed at l,650 F. A 30,000 p.s.i. stress level was found to be one in which sigma phase would form in a realistic time. If the sigma phase was excessive, it would cause premature failure. The following Table III compares typical alloys of Table 11 under these stress rupture conditions.

TABLE III Stress Rupture Properties (As Cast) 1,650 F.30,000 p.s.i.

From Tables 11 and 111, it is easily seen that alloy Examples 1 and 2, having Ti in the range of 3.5 4.3% and a phase computation number N less than 2.47, have significantly greater stress rupture life than the other alloy examples, all of which lie outside the range of the present invention. Photomicrographic studies of alloys 1, 2 and 12 indicated the general absence of sigma phase and hence a different kind of alloy structure than that of the other alloys which showed abundant amounts of sigma phase. Because of these two different kinds of phase structures, the stress rupture life of the alloy within the scope of the present invention at 4.0-4.3% Ti is significantly greater and, in addition, the alloy is more stable. The tendency toward lower strength, as was mentioned before in connection with lower Ti, is shown by a comparison of the data for alloys 1 and 2, with alloy 2 being at 3.5% Ti.

The following Table IV lists the tensile properties of some of the alloys of Table 11. These data show that while the stress rupture life of the alloy of the present invention is significantly increased, there is substantially no change in tensile strength. However, there is an increase in ductility.

The testing reported in Tables III and IV was performed on unmachined, cast-to-size bars. In Table IV, the terms UTS means Ultimate Tensile Strength in thousands of pounds per square inch and the El means Elongation in inches per l-inch gage length.

The above Tables III and IV clearly show that maintaining the titanium level at or below about 4.3 weight percent and the phase computation or electron vacancy number N at a maximum of 2.47 is extremely beneficial in providing long time stability. As has been indicated before, this is based on the substantial elimination of the formation of sigma phase. Alloy example 1 of Table II with 4.2% Ti and N,.;, of 2.33, was found to be free of the acicular phase after 1,000 hours at l,650 F. Alloy 1 has about 3 or more times the stress rupture life of alloys 9, 10 or 11, which have an N of 2.68 or higher with Ti in the range of 4.5-5.4% and thus are outside the range of the present invention. As described before, the alloy of the present invention defines the Ti range as 4.04.3% along with an N of 2.47 or less.

Although alloy 5 contains titanium within the range of the present invention and chromium at about 10 weight percent, such alloy contains molybdenum in a range of greater than that of the present invention. The increased Mo detracts significantly from the stress rupture properties of alloy 5 as shown by the data of Table III. The inclusion of molybdenum at greater than about 4 weight percent can greatly decrease the ductility of the type of alloy to which the present invention relates. For example, alloys 3, 4 and 5 have significantly lower tensile ductility as shown in Table IV. In addition, the stress rupture ductility was low for alloy 5 as shown in Table III.

As was mentioned above, increased chromium level promotes sigma phase formation and loss of ductility. It has been found that lower chromium content suppresses sigma phase formation, increases ductility and increases overall strength without notable decrease in oxidation resistance. If the chromium content is too high as shown by alloy 8, including about 14 weight percent Cr, even the reduction of titanium to about 4 weight percent does not sufficiently lower the N and hence does not suppress the sigma phase formation. Alloy 8 shows an extreme abundance of sigma phase after exposure and was weak in 1,850 F. rupture, failing in 3.7 hours. It failed to show sufficient ductility after 500 and 1,000 hours exposure at l,650 F. to warrant its stress rupture evaluation at l,650 F .30,000 p.s.i. Because the physical property tests were conducted in air under the stated conditions, it is to be noted that lowering the Cr content does not decrease the oxidation resistance since stress rupture life also reflects any strength deterioration due to oxidation.

With regard to cobalt, alloys 6 and 7 were designed to show the effect of cobalt on the properties of the type of alloy to which the present invention relates. The mechanical properties generally followed those of similar alloys in the Tables because of the formation of sigma phase. Even with low Cr.

sigma phase was formed because of the higher Ti content at 5.3% [t is concluded that cobalt alone between about 10-20 weight percent does not eliminate the sigma phase formation without concurrently reducing the Ti level.

Testing and photomicrographic studies of alloy 12 of 'lablc 11 showed the alloy to be free of sigma phase at 4.3% Ti after 5,000 hours at 1,550 F. At 1,800 F. and 10.000 p.s.i. stress, a specimen of alloy 12 had a life of 5,532 hours.

Thus although there is evidence of sigma phase formation in alloys at 4.5% Ti, otherwise within the scope of this invention, alloys at about 4.3% are sigma free.

The following Table V lists the compositions and phase computation numbers (N,,;,) for a series of known alloys. Although the numerical values for the compositions of alloys D through K of Table V appear to be similar to that of the present invention, it is to be noted that the calculated values for their phase computation number (N,,;,) are all well in excess of the 2.47 definition for the alloy of the present invention. As shown from the previous data, all of the alloys in Table V can be classified as sigma-prone alloys and therefore are limited to a long time stress rupture life at elevated temperatures less than what might be expected without consideration of the presence and detrimental effect of sigma phase.

TABLE V Compositions in Weight Base: 0.18 C; 0.02 B; 0.05 Zr, Bal. Ni and Incidental Impurities Although the present invention has been described in connection with specific examples, it will be understood by those skilled in the art of metallurgy the modifications and variations of which this invention is capable.

What is claimed is:

1. A nickel base alloy of improved stress rupture stability, having a phase computation number (N at a maximum of 2.47 and consisting essentially of, by weight, 0.10.2% C, 8-11% Cr, 4.04.3% Ti, 5-6% l0-0.03-0.09% Zr, 10-20% Co, 24% Mo, 0.7-1.2% V, 0.010.02% B, with the balance nickel and incidental impurities, the alloy structure characterized by the absence of sigma phase.

2. The alloy of claim 1 in which the Cr is 9-1 1%, the Co is 14-16%.

3. The alloy of claim 2 in which the C is 0. 15-02%, the Cr is 910% and the Al is 5.35.7%.

232 33 UNITED STATES. PATENT OFFICE v CERTIFICATE OF CORREGTION Patent No. 3, 642,469 Dated February 15, 1972 Inventofls) Earl Warren Ross and Howard Thomas McHenry It is certified that error appears in theabove-identified patent and that said Letters Patent are hereby corrected as shown below:

Column 8, line 24, change "5 6% 10-" to 5 6% Al,

Signed and sealed this 20th day of June 1972.

(SEAL) Attest:

ROBERT GOT'I'SCHALK Commissioner of Patents EDWARD M.FLETCHER,JR. Attesting Officer 

2. The alloy of claim 1 in which the Cr is 9- 11%, the Co is 14-16%.
 3. The alloy of claim 2 in which the C is 0.15- 0.2%, the Cr is 9- 10% and the Al is 5.3- 5.7%. 